Because of their excellent fluidity and castability, eutectic and hypoeutectic aluminum-silicon alloys are widely used in the production of aluminum castings. In an unmodified state, the eutectic silicon phase is present as coarse plates with sharp sides and ends often referred to as acicular silicon. The presence of acicular silicon results in castings which have low percent elongation, low impact properties and poor machinability.
Strontium has been shown to be effective in refining or modifying coarse acicular silicon into a fine, interconnected fibrous structure. In general, small quantities of strontium between 100 to 200 ppm are sufficient to produce a fine, fibrous eutectic silicon which in turn significantly improves the mechanical properties and machining characteristics of the aluminum casting. U.S. Pat. No. 3,466,170 issued Sep. 9, 1969 to Dunkel et al. recognizes the benefit of adding strontium either as a pure metal or as an AlSr alloy with 7 net percent Sr.
Because strontium metal is very reactive with oxygen, nitrogen and moisture, its use as a modifying agent is limited. In most cases, strontium is added in the form of a master alloy.
The publication by Pekguleryuz et al., "Conditions for strontium master alloy addition to A356 melts", Trans. Am. Foundrymen's Soc. (1989) considers the use of a 55 wt. % Sr 45 wt. % Al alloy as a master alloy to modify aluminum-silicon alloys. This alloy largely comprises the intermetallics Al4Sr and Al2Sr. This document does not, however, disclose the alloy being in the form of granules or powder.
U.S. Pat. No. 3,567,429 issued Mar. 2, 1971 to Dunkel et al. teaches the use of a strontium silicon-aluminum master alloy which has a strontium content higher than 7%. Strontium-silicon aluminum master alloys are no longer widely used for modifying aluminum-silicon casting alloys, since in most cases the strontium is present as a high melting temperature intermetallic phase such as Al.sub.2 Sr.sub.2 Si or SrSi.sub.2 which dissolves very slowly at molten aluminum processing temperatures, typically 760.degree. C. or lower. As reported by John E. Gruzleski and Bernard M. Closset ("The Treatment of Liquid Aluminum-Silicon Alloys", American Foundrymen's Society Inc., 1990, pages 31-39), a 10% strontium-aluminum binary master alloy dissolves twice as fast in a A356 aluminum-silicon casting alloy as a 10% strontium-14% silicon-aluminum ternary master alloy at all melt temperatures ranging between 670 to 775.degree. C. Similar results are found in U.S. Pat. No. 5,045,110, issued Sep. 3, 1991 to Vader et al., reporting dissolution times between 20 and 30 minutes for 10% strontium-14% silicon-aluminum master alloys in ingot form. In contrast, U.S. Pat. No. 4,576,791, discussed below, teaches that 5-10% strontium-aluminum binary alloys in rod form and which contain titanium and boron grain refiners dissolve in 1 minute. In addition, the customary process used to produce strontium-silicon master alloys results in substantial quantities of detrimental impurities including iron, barium and calcium often being present in the master alloy.
U.S. Pat. No. 4,108,646 teaches the use of a master composition consisting of strontium-silicon in particulate form pressed into a briquette with aluminum or aluminum-silicon particles. The briquettes, having a master composition of between 3 to 37% strontium by weight, are then added to an aluminum-silicon casting alloy to modify its structure. This master composition is less efficient than aluminum-strontium binary master alloys since the strontium is present as SrSi.sub.2 particles which, as discussed above, dissolve slowly and contain detrimental impurities including up to 4% iron and 1 to 3% calcium.
Aluminum-strontium binary alloys are now widely used for modifying aluminum castings; however, it has been difficult to increase the strontium content of these binary master alloys. This is best explained in the context of the aluminum-strontium binary equilibrium phase diagram of FIG. 1. The phase diagram contains two low melting point eutectics, one at about 3.5% strontium, the second at 90% strontium. On the aluminum rich side, the eutectic containing alloys range from about 0% to 44% strontium. On the strontium rich side, the eutectic containing alloys range from about 77% to 100% strontium. In the final solidified state, these eutectic alloys contain in varying proportions a eutectic phase which is very finely divided and melts at low temperatures, 654.degree. C. in the case of the aluminum rich eutectic and 580.degree. C. for the strontium rich eutectic. These finely divided eutectic phases are more ductile and dissolve more rapidly than the higher melting point intermetallic alloy phases which are present between about 44% to 77% strontium. Since these intermetallic alloys contain no low melting point, finely divided eutectic phase, they are more brittle and dissolve much more slowly than the eutectic containing alloys. The presence of these high melting point intermetallics alloys has placed a significant limitation on the amount of strontium which can be effectively contained in commercial aluminum-strontium binary master alloys. In this specification, the term "intermetallic alloys" denotes alloys containing between approximately 40% to 81% strontium by weight. These alloys are dominated by the Al.sub.4 Sr, Al.sub.2 Sr and AlSr intermetallics and contain only minimal or no eutectic phase.
As discussed in "Phase Diagrams for Ceramists" compiled by the National Bureau of Standards, published by The American Ceramic Society Inc., Volume 1, pages 9-14, FIG. 1 as a binary equilibrium phase diagram shows the relationships between composition and temperature assuming all phases are in equilibrium with each other. These compositional relationships are only valid if the rate of solidification is slow enough to allow the phases to reach compositional equilibrium at every instant. A more rapid rate of solidification will lead to quite different compositional results.
As shown in FIG. 1, when a liquid alloy containing 10 % strontium is cooled, solidification begins at about 815.degree. C. The first solid phase to precipitate is primary Al.sub.4 Sr intermetallic which contains approximately 44% strontium. As the melt temperature continues to decrease during solidification, more and more of this primary Al.sub.4 Sr intermetallic phase precipitates. The primary Al.sub.4 Sr intermetallic phase is present as massive interconnected plates or needles which are shown two-dimensionally in the photomicrograph given in FIG. 2. A three-dimensional view of the interconnected network of primary Al.sub.4 Sr plates is shown by FIG. 3 taken using a stereomicroscope.
When the melt temperature cools to 654.degree. C., the primary Al.sub.4 Sr intermetallic phase stops precipitating and the remaining amount of liquid alloy solidifies as a very finely divided, ductile eutectic phase. The eutectic phase is shown in FIG. 2 by the light regions surrounding the large primary Al.sub.4 Sr needles. The eutectic phase is much more finely divided than the Al.sub.4 Sr intermetallic phase as evidenced by the lack of resolution of the eutectic phase at 50 times magnification.
The quantity of primary intermetallic Al.sub.4 Sr phase present in the final solidified alloy will depend on the rate at which freezing took place between 815.degree. C. to 654.degree. C. If the alloy were allowed to freeze very slowly so that equilibrium is achieved at each instant of cooling, then the quantity of primary Al.sub.4 Sr intermetallic phase in the final alloy will be given from the equilibrium phase diagram in FIG. 1 using the lever rule, that is for a 10% strontium alloy ##EQU1##
As discussed in "Phase Diagrams for Ceramists" above, a more rapid rate of solidification which does not allow phase equilibrium at each instant will lead to quite different compositional results.
A more rapidly solidified alloy will contain less than 16% primary Al.sub.4 Sr intermetallic phase with the quantity of primary Al.sub.4 Sr decreasing as the rate of freezing increases. This reduction in the quantity of the primary Al.sub.4 Sr intermetallic phase as the rate of solidification increases is due to the shorter period of time spent by the freezing alloy in the 815.degree. C. to 654.degree. C. temperature range where the primary Al.sub.4 Sr precipitates. Hence rapid solidification leads to less primary intermetallic phase and correspondingly an increase in the quantity of eutectic phase in the final solidified alloy. For a 10% strontium-90% aluminum master alloy, the maximum quantity of primary Al.sub.4 Sr intermetallic phase is 16%, and correspondingly the minimum quantity of eutectic phase is 84%, which occurs when cooling rates are slow enough to allow for phase equilibria.
U.S. Pat. No. 4,576,791 states that a 10% strontium-aluminum alloy rod, which contains a maximum of only 16% primary Al.sub.4 Sr intermetallic phase and at the very minimum 84% finely divided eutectic phase, normally dissolves so slowly as to be unsuitable for use as master alloy in rod form. This is due to the presence of relatively large crystals of Al.sub.4 Sr primary intermetallic phase ranging from 5 to 300 microns as viewed two-dimensionally through a microscope. The patentee meets this problem by providing 0.2 to 5% titanium and up to 1% boron in the master alloy to refine the typical Al.sub.4 Sr primary intermetallic two-dimensional crystal size to 20 to 100 microns. Reducing the size of the Al.sub.4 Sr primary intermetallics increases the ductility of the rod thereby enabling it to be coiled and uncoiled during feeding and also shortens the dissolution time to approximately 1 minute which is required for launder additions. The addition of titanium and boron enables strontium concentrations in the master alloy to be increased to 20% Sr by weight, in the preferred embodiment to 10% Sr. Refining the size of the primary Al.sub.4 Sr intermetallic phase is effective up to a maximum of 20% strontium beyond which the alloys are unsuitable for use in rod form.
In U.S. Pat. No. 4,576,791, the Al.sub.4 Sr primary phase crystals are referred to as ranging from 5 to 300 microns in size. It is important to note, however, that this size description may be misleading since it is based on a two-dimensional microscopic view of a polished sample (FIG. 2). In actuality the primary intermetallic phase forms first during solidification as a three-dimensional network of crystals. Even though in a two-dimensional microscopic view the Al.sub.4 Sr intermetallics appear as discrete needles sized less than 300 microns, in actuality these intermetallic crystals form an interconnecting network of plates surrounded by very finely divided eutectic phase which is the last phase to solidify. FIG. 3 shows the three-dimensional interconnected plates of Al.sub.4 Sr primary intermetallic phase present in a 10% Sr-90% Al alloy. The amount of three-dimensional interconnection increases as the strontium concentration increases in the alloy. Hence, in the prior art there has been an upper strontium concentration limit. Beyond this upper strontium concentration limit, the three-dimensional network of interconnected primary phase intermetallic crystals becomes too large and the quantity of finely divided eutectic surrounding the plates too small rendering the alloy unusable due to the slow dissolution and brittleness of these large intermetallic networks.
A different approach to the problem caused by Al.sub.4 Sr plates is found in U.S. Pat. Nos. 5,045,110 and 5,205,986, issued Sep. 3, 1991 and Apr. 27, 1993, respectively, in the name of Shell Research Ltd. These patents teach that the strontium concentration in binary aluminum rich-strontium master alloys can be increased to 30% or 35% Sr by weight by further refining the grain size and reducing the quantity of the Al.sub.4 Sr primary intermetallic phase as a result of atomizing the liquid alloy at very rapid cooling rates of 10.sup.2 to 10.sup.4.degree. C. sec. By this process both the quantity and the size of the primary Al.sub.4 Sr intermetallic phase which precipitates first is reduced and the quantity of finely divided, more ductile eutectic phase is increased proportionately.
FIG. 4 is a photomicrograph taken at 500 times magnification of a 10% strontium-90% aluminum alloy rod produced from a rapidly solidified atomized alloy as in U.S. Pat. Nos. 5,045,110 and 5,205,986. When compared to FIG. 2 which is a photomicrograph taken at only 50 times magnification (10 times lower magnification than FIG. 4) of a 10% strontium-90% aluminum alloy cast in a permanent mould at moderate solidification rates, it is evident that the rapid solidification rates resulting during atomization greatly reduces the size and quantity of the primary Al.sub.4 Sr intermetallic phase. Titanium and boron may also be added to the master alloy to further refine the structure. By reducing the quantity and refining the size of the primary Al.sub.4 Sr intermetallic phase and also increasing the quantity of very finely divided, ductile eutectic phase, the patents teach that the strontium concentration in aluminum-strontium master alloys can be increased up to 35% Sr by weight. The atomized solid particles, each of which contains both a finely divided Al.sub.4 Sr intermetallic phase and a eutectic phase, are consolidated by an extrusion process into a rod for in-line addition to a launder, this rod having "sufficient ductility to enable coiling and decoiling".
Although as detailed by Gruzleski and Closset above, a 90% strontium rich-aluminum master alloy is also available but is of limited use as a master alloy. This strontium rich master alloy consists of 100% finely divided eutectic phase with no intermetallic phases present and has very limited application since it can only be used when the aluminum-silicon casting alloy melt temperature is below about 720.degree. C. When added to an aluminum alloy melt, the 90% strontium alloy first melts and the 90% strontium enriched liquid then dissolves to dilute levels of 150 to 200 ppm Sr. During this dissolution, the local liquid composition must become diluted from 90% strontium down to less than 0.02% Sr (150-200 ppm Sr). During this dilution, the local melt composition must pass through the range of high melting point intermetallic alloy compositions from 77% to 44% strontium and these intermetallic phases will precipitate during dissolution as solid intermetallic phases which stop or further retard strontium dissolution. At melt temperatures below 720.degree. C., the 90% strontium alloy dissolves exothermically releasing sufficient heat to raise the aluminum-silicon alloy melt temperature locally to a sufficiently high level as to avoid the formation of the high melting Al.sub.4 Sr and Al.sub.2 Sr intermetallic phases. Hence at melt temperatures below 720.degree. C., the 90% Sr-10% Al alloy dissolves rapidly with high recovery. At melt temperatures above about 720.degree. C., this exothermic reaction diminishes and insufficient heat is generated. This results in the formation of the Al.sub.2 Sr and Al.sub.4 Sr intermetallic phases during dissolution. The presence of the high melting Al.sub.4 Sr and Al.sub.2 Sr intermetallic phases effectively retards dissolution and results in poor strontium recovery.
Thus, as taught by the prior art discussed above, the presence of high melting point primary intermetallic phases between 44% and 77% strontium by weight has placed significant limitations on the use of aluminum-strontium master alloys.
Hitherto, the useful aluminum-strontium master alloys have been alloys which contain substantial quantities of very finely divided, ductile, low melting point eutectic phase. In the case of aluminum rich-strontium master alloys ranging from 5 to 35% strontium, the alloy consists of a mixture of primary Al.sub.4 Sr intermetallic phases surrounded by finely divided eutectic phase. The primary Al.sub.4 Sr intermetallic phase is present as a three-dimensional network of interconnected plates which under normal solidification rates can be quite coarse in size. All of the prior art teaches that the only method of increasing the strontium concentration in these aluminum rich master alloys while allowing acceptable rates of dissolution of the alloys in molten aluminum is to minimize the quantity and refine the size of the interconnected network of Al.sub.4 Sr plates and to maximize the quantity of the more ductile, very fine eutectic phase.
A separate use of the aluminum-strontium alloy of this invention is as an inoculant for cast iron.
Inoculation is a process in which formation of metastable carbides is suppressed in cast iron. Instead, graphite which is the equilibrium phase is allowed to form.
Industrial solidification rates which are usually between 0.1 to 10.degree. C./sec do not normally allow the formation of graphite at thinner sections of the castings. Inoculation provides substrates or nuclei for the nucleation of graphite. These substrates or nucleation sites are believed to be sulfides. Strong sulfide formers such as calcium and strontium are added to molten iron for inoculation. A good source of information on the possible nucleation mechanisms and practice of inoculating cast iron can be found in "The Modern Inoculating Practices for Gray and Ductile Iron", Proceedings of AFS-CMI Conference, Feb. 6-7, 1979.
U.S. Pat. No. 4,666,516, issued May 19, 1987 to Hornung et al., also discusses the manner in which the form taken by carbon present in cast iron greatly affects its characteristics. If in the form of iron carbide (known as "chill"), the cast iron is brittle ("white cast iron") but if in the form of flake graphite the cast iron is soft and machinable ("gray cast iron"). The spherical form of graphite produces higher strength and improved ductility ("ductile cast iron").
Ferrosilicon has been used as an inoculant to promote the formation of graphite particularly in nodular or spherical form and U.S. Pat. No. 3,527,597, issued Sep. 8, 1970 to Dawson et al. teaches that an important inoculant can be obtained by including 0.1 to 10% strontium metal and maintaining a low calcium content. Commercial grade ferrosilicon contains calcium as an impurity. As pointed out in the patent: "pure ferrosilicon has very little inoculating effect when added to cast iron; commercial foundry grade ferrosilicon depends upon its content of small amounts of minor elements, notably aluminum and calcium, for stimulating the inoculating effect".
It has now been found that by using the strontium-aluminum alloy of this invention a further improved inoculant effect can be obtained. It has been found that the most useful composition is ferrosilicon together with the 80% Sr alloy. In the case of gray cast iron, chill was eliminated and the amount of type D graphite minimized by the combined addition. In the case of the ductile iron the combined addition minimized the amount of chill and increased the number of nodules.
In U.S. Pat. No. 3,527,597 metallic strontium was added to gray iron along with FeSi. The required amount of strontium was much higher than if it had been alloyed with FeSi and Si. In the case of alloying with silicon only a maximum of 65% Sr could be obtained in the final alloy due to the nature of the dissolution reaction. In the present invention, the additive can contain up to 80% strontium and only a minute amount is required for successful inoculation. U.S. Pat. No. 3,527,597 teaches that the only FeSi grade that can be alloyed with Sr as a potent inoculating agent is a low calcium FeSi which normally contains less then 0.35% Ca. Above this Ca concentration the potency of the inoculant is diminished. The patent also discloses the separate addition of SiSr alloy alone or together with FeSi containing low and normal concentration of calcium. SiSr with low calcium FeSi produced a significant inoculant effect on gray cast iron.